DEVELOPMENT OF ROLLING TEXTURES IN AN AUSTENITIC STAINLESS STEEL

Three dimensional texture analysis by means of orientation distribution functions (ODF) was used to examine the texture development during rolling at 473 K in an austenitic stainless steel. With the help of ODFs results, the different stages of texture development could be assigned to the existing theories of heterogeneous deformation mechanisms of low SFE face-centred cubic metals. The texture at very low degree of rolling consists of two limited orientation tubes with their fibre axes (110)[]ND and (ll0)60ND and agrees with the predictions made by Taylor model. With further deformation, twinning causes the reduction of ={112} (111) component and leads to the formation of twin {552}(115). Abnormal slip on slip planes parallel to the twin boundaries rotates the twins into the {332} (113 and 111 (110) positions. The shear bands formation in the rotated twin-matrix lamellae changes their orientations near to {011} (100) and {011} (112) positions. Finally, normal slip again continues and sharpens the brass-type rolling texture.


INTRODUCTION
The study of the development of rolling textures in cubic metals in general and austenitic stainless steel in particular at ambient temperature, is of great scientific and technological importance, since these control further processing stages and thereby the resulting properties (i.e.Young's modulus or plastic anisotropy).It also gives more information on the active deformation mechanisms in polycrystalline materials.
There is a general consent in the literatures (Wassermann and Grewen, 1962;  Liicke, 1981; Bacroix and Jonas, 1988; Hirsch and Liicke, 1988) that face-centred cubic (FCC) metals and alloys possess different rolling texture (i.e.copper-type and brass-type).The transition texture in between copper-type and brass-type is also perceptible.The possible causes for different rolling textures are mainly rolling degree, rolling temperature, starting texture, grain shape and finally the stacking fault energy (SFE).As the SFE increases the texture transition from brass to copper-type takes place and has been systematically surveyed as a function of solute element (Zn) in copper (Alam, Mengelberg and Liicke, 1967)  and as a function of rolling temperature in austenitic stainless steel (Goodman and Hu, 1964), copper (Hu and Goodman, 1963) and silver (Hu and Cline, 1961;  Hu, Cline and Goodman, 1961).The FCC rolling textures may be described either in terms of ideal orientations (Hu, Sperry and Beck, 1952) or in terms of complete or limited fibre axis (Grewen and Wassermann, 1955; Wassermann,  1963).According to the concept of ideal orientations, the copper-type texture is characterised mainly by the orientations {112}(111), {123}(634) and {011}(112), in doing so the indexing of {123}(634) component can adopt somewhat { 123} (412) or { 124} (211), whereas the brass-type texture is domin- ated by {011}(112) and to a lesser extent by {011}(100).The transition from copper to brass-type textures is described by the disappearance of { 112} (111) component.On the contrary, the copper-type texture has been described by taking into consideration not only the ideal orientations {112}(111) and {011}(112) but also the spread ranges extended between them with the aid of two limited (111) fibres formed by rotations upto +30 about certain {111} poles of two ideal components.Out of the two fibre axes, one is taken under 19 to RD on the periphery of the {111} pole figure and the other under 19 from ND on the ND/RD radius of the pole figure.
Extensive metallographic studies of the deformation structure by optical and transmission electron microscopy (TEM) have been performed on FCC metals like copper and brass and these microscopical results lead to a number of new aspects in the interpretation of deformation structure and texture development (Duggan et al., 1978; Hatherly, 1978; Malin and Hatherly, 1979; Wakefield and  Hatherly, 1981).According to Malin and Hatherly (1979) the deformation sequence in high and medium SFE metals (i.e.copper) involves the formation of equiaxed cells, microbands, clustering of microbands and shear bands.In these metals copper-type rolling texture forms which can be determined theoretically from the calculations based on Taylor model (Dillamore, Butler and Green, 1968;  Bunge, 1970; Dillamore and Katoh, 1971; Dillamore and Katoh, 1974), Taylor model adapted to both slip and twinning (Van Houtte and Aernoudt, 1976; Van  Houtte, 1978a; Van Houtte, 1978b) and relaxed constraints model (Honneff and  Mecking, 1981; Van Houtte, 1981).With decrease in SFE, an increasing tendency of mechanical twinning during deformation is often considered as the mechanism (Wassermann, 1963) which causes the change from the copper-type to the brass-type rolling texture.In metals of low SFE such as 70-30 brass (Fargette and Whitwham, 1976; Grewen, Noda and Sauer, 1977; Duggan et al., 1978;  Hutchinson, Duggan and Hatherly, 1979) and stainless steel (Blicharski and  Gorczyca, 1978), deformation begins by the glide of partial dislocations, followed by mechanical twinning, the rotation of the twin-parent lamellae into alignment with the rolling plane, the development of shear bands and the resumption of normal octahedral slip in the recovered crystallites of the shear bands.But the effects of the mechanical twinning on further deformation and on the formation of rolling textures in austenitic stainless steel, which become even more complicated by shear band formation, are not yet completely understood.Furthermore, the characterisation of the rolling textures developed in austenitic stainless steel above ambient temperature was based on the pole figure analysis (Goodman and Hu, 1964) which possesses very limited resolving power in revealing the differences which are more of quantitative than of qualitative nature.
The aim of the present investigation was to clarify the microscopical observa- tions of orientations and their statistical appearance in the textures of austenitic stainless steel cold rolled at 473 K by means of the three-dimensional ODF analysis which describes the texture in a more complete, exact and explicit fashion.The results also outline the process of texture development during cold rolling.

EXPERIMENTAL PROCEDURE
A commercially produced hot band (HB) of an austenitic stainless steel of nominal composition (by wt%): 19.25% Cr, 8.40% Ni, 0.05% C, 1.24% Mn, 0.48% Si, 0.02% P, 0.02% S was used for the present investigation.As per supplier's information 2.95 mm thick HB strip was finished at about 1173 K and coiled at about 1023 K.In order to remove any oxides and scales from the surfaces, the hot band was pickled in a pickling solution (by volume %) of 10% HNO3 + 2% HF + 88% distilled water at a temperature 323 K for a period of 5 minutes and then rinsed with water and dried.
The hot band was cold rolled (CR) at a temperature 473 K without lubrication.
The cold rolling operation was oriented parallel to the original hot rolling direction and carried out in such a way that the true strain per pass was about 0.05.In the interval between the passes specimens were kept in the furnace at 473 K for a period of 5 minutes in order to maintain the temperature.The two-high mill used had 55 mm dia rollers.The cumulative rolling reductions investigated were from 10 to 90% at intervals of 10%.
Optical metallography was performed on specimen of HB, which was mechani- cally polished and electrolytically etched in saturated oxalic acid solution at 10 volts for 2 minutes, at section (S 0) cut perpendicular to the transverse direction containing (R) and normal direction (N) RN.Here S =0 denotes the mid- thickness of the specimen.
X-ray texture measurements were performed for HB rolled at 0, 30, 50, 70 and 90% reduction in thickness on an automatic texture goniometer, using MoK radiation.For each ground and etched specimen (20 mm x 14 mm), the texture was determined at the section (S 0) parallel to the rolling plane by measuring four incomplete pole figures of the plane {111}, {200}, {220} and {113}, using Schulz back reflection technique (Schulz, 1949).The diffracted intensity was recorded continuously every 5 along concentric circles in the angular range from 0 to 75 in steps of 5.The measured intensity was subjected to background, geometrical and defocussing corrections using a random specimen of pressed and sintered austenitic stainless steel powders.Each ODF was calculated from the data of four incomplete pole figures following the series expansion method of Bunge (1982) and using the pseudo-normalisation technique of Kern and  Bergmann (1978).The series were extended to the degree 22.The texture strength as expressed by the texture index J (Truszkowski et al., 1978) which defines the mean square deviation of the ODF from random distribution, was also determined in each case.The inverse pole figure measurements were also made for the cold rolled specimens parallel to the rolling plane.For these, the measured integrated intensities of selected reflections were used for determining the pole densities of austenite grains having their plane normals parallel to normal direction (ND) of the sample (Singh, 1989) and the pole density of austenite grains was expressed as lh,l/Rht,! 1 E lh,l/Rhk tn where /, R and rn represent the integrated intensity, theoretically calculated factor and number of reflections respectively.

EXPERIMENTAL RESULTS
The industrially produced hot band of austenitic stainless steel was fully austenitic.The optical micrograph of the HB in the RN section near the centre (S-0) is shown in Figure 1 and indicates mostly the elongated grains.Some recrystallised grains are also present.The dark bands in the micrograph appear to be thin sheet like bands of heavily deformed austenite grains.The average grain size measured on the section (S 0) parallel to the rolling plane was about 7 #m.Figure 2 shows the three dimensional orientation density distribution (i.e. ODF) of crystallites for the HB in constant P2 sections through Euler space.This ODF exhibits maxima of orientations which are given in Table 1 along with their Euler angles (ql, tp, tP2).The main features of the texture components present in this ODF are rolling (copper-type) and recrystallisation textures.Singh, Ramas- wamy and Suryanarayana (1991) have discussed the texture evolution in the present steel during hot rolling in detail and characterised the rolling texture components, which were retained, as {011 } ( 112), (123} (634) and { 112} ( 111 ) orientations and also described the orientations {001}(100), RD rotated cubes {013} (100), {012} (100) and Goss (011) (100) as the recrystallised components.
Upon rolling the HB at 473 K, very limited amount of cr'-martensite (about 16.0 vo|.% at 90% reduction) was produced at the centre, which was not sufficient for its texture measurements (Singh, 1989).
The ND inverse pole figure densities of the selected planes [i.e.(lll)r, (200)r, (220) and (311)] parallel to the rolling plane of the cold rolled specimens, with increasing rolling reductions (range: 10-90%) are shown in Figure 3. Although, the pole density of the plane (220)r is relatively higher than that of the other planes, it increases rapidly between the reductions 10 to 50%, and thereafter, remains more or less constant upto 80% reduction and then there is an increasing trend after 80% reduction.In contrast, the pole density of (111) remains relatively constant upto 40% reduction, and thereafter, increases to a maximum value at 80% reduction and then, there is a decreasing trend after 80% reduction.
The pole densities of other two planes (311), and (200), are continuously decreasing with increasing reduction.
increasing deformation and also there is a decrease in the density of the cube orientation with deformation.

DISCUSSION
The microstructure [Figure (1)] at the centre of the HB reveals the presence of deformed as well as recrystallised grains.Further, the interior layers of the HB retain considerably a deformation structure.The ODF analysis of the HB also indicates the presence of textural elements of both the retained rolling textures (i.e.copper-type) and recrystallisation texture (i.e.cube and RD rotated cubes).
The results of the ODFs analysis of the deformed austenite at the centre of the HB cold rolled at 473 K show that at low degrees of rolling (<50%), the observed texture components are distributed along the two limited orientation tubes.The first tube with its fibre axis (ll0)//ND stretches along {Pl at I#2 0 and tp 45 and includes the orientations {011}110) at pl=b=P2=0, 45 and 0 and {011}(112) at tPl tp tp2 35, 45 and 0. The other tube inclined to constant 2 sections runs through the Euler space from (225} (554) ({112} (111)) at tp, 2=90, 30 and 45 over {123)(634) at ql, tp, tP2=59, 37 and 63 , to (011}(112) at tPl, tp, tP2=35, 45 and 90 and has its fibre axis (110) inclined ---60 from ND towards RD.The orientation element {011}( 112) is at the intersection point of the above two limited tubes and has the highest density amongst all the texture components.The exceptional high density of {011}( 112) component is partly due to the effect of starting texture of the HB (Figure 2) which shows the highest density of the orientation {011}(112) and partly due to the flow of orientations of the scattered zone between the cube components {001}(100) at Pl, {P, 1#2 =0, 0 and 0 and Goss component {011}(100) at {P2 0, 45 and 0 , along the line parallel to tp at {Pl--" lP2 0 from the cube to Goss component (i.e.metastable orientation) and then along the line parallel to l at p=45 and tP2=0 , from Goss to the component (011}(112) (i.e.stable orientation) at Pl, tp, tP2-35, 45 and 0 , obtained as a result of rolling.The orientation {225}(554) at which the orientation tube with (110) 60ND fibre starts, differs by 5 from the orientation (112}(111) at tPl, tp, 2 90, 35 and 45 and by 3 from the theoretically calculated orientation {4 4 11} (11 11 8) at P2=90, 27 and 45 by Dillamore, Butler and Green (1968) for Taylor type deformation.An increase in orientation densities for all texture components of the two limited tubes i.e. (ll0)//ND and (ll0)60ND (Figures 5 and 6) occurs with increasing strain (e.g.<30% reduction).All these evidences suggest that the development of rolling textures at low degree of rolling (<50%) agrees fairly well with the predictions of Taylor model (Dillamore, Butler and Green, 1968).At 50% rolling reduction, the decrease in the otherwise stable {112}(111) orientation marks the beginning of mechanical twinning (Wassermann, 1963) whereby the orientation (112} (111) undergoes mechanical twinning during rolling to form {552} (115) at , q, 2 90, 75 and 45.Wassermann (1963) has indicated that the twinning is the most favourable deformation process which causes the beginning of transition of the rolling texture from copper to brass-type.
It has been reported by TEM studies of the microstructures in rolled polycrystalline 70-30 brass (Duggan et al., 1978) and silver (Pospiech et al., 1975)   that twinning becomes an important deformation mechanism at certain strain level but instead of normal rotation of the twin component (552( 115) towards (011)(100), an abnormal rotation of the twin-matrix lamellae towards (111) (112) position occurs.The orientation (111) ( 112) is a position of the twin planes parallel to the rolling plane.The results of ODF analysis (Figure 7) where the twin orientation {552)(115) at intermediate degrees of rolling (>50%), shifts to lower angles instead of rotating towards (011}(100) position, prove that a rotation towards (111} (112) position takes place, whereas the orientation density of the orientation (011)(100) remains more or less constant (i.e.rather increases very slightly).
This abnormal rotation towards (111} ( 112) is mainly due to the preferred slip on slip planes parallel to the coherent twin boundaries of the twin lamellae.In metals of low SFE, the normal slip on the slip planes of highest critical resolved shear stress, is severely limited mainly because of the very small width of the twin lamellae (Peissker, 1965).Further, the shear stress for the abnormal slip starts decreasing and finally becomes zero as the twin-parent lamellae become oriented with the rolling plane, and ultimately the abnormal rotation stops and, as detected in Figure 7, the statistical maxima of the rotated orientations is observed near {332)(113), which is 10 away from {111)(112) position.The rotation of the matrix orientation { 112} ( 111 ) towards { 111} (112) could not be observed in the ODF results (Figure 7) mainly because the obtained ODF is the reduced ODF.However, Hirsch, Virnich and Liicke (1981) have reported the rotation of matrix orientation towards {111} (112) position in 70-30 brass with the help of complete ODF.
Besides the component {332}(113) at intermediate degree of rolling, an orientation near {111}(110), also with {111} plane nearly parallel to the rolling plane appears.Since {111} (110) develops simultaneously with the {332} (113), it can be assumed that this position is formed by the preferential slip on one of the two active slip systems parallel to the twin boundary that rotates the twin component {552}(115) and other possible twin orientation of {123}(634) component.The twin of { 123} (634) component is symmetrically equivalent to {123} (634) position.This causes the rotation of the preferred slip direction parallel to the rolling plane.
Since the rotated twin-matrix lamellae finally no longer contribute to deforma- tion process by slip, on further rolling shear bands form in the highly unstable structure of aligned twin-matrix lamellae.Formation of shear bands has been observed by many investigators (Duggan et al., 1978; Hatherly, 1978; Blicharski  and Gorczyca, 1978; Wakefield and Hatherly, 1981) in their TEM studies.
Further, the cyclic variation in the orientation density of {332} (113) component after 50% reduction indicates a competition between mechanical twinning and the associated abnormal slip on one hand and shear bands formation in the rotated twin-matrix lamellal on the other.Twinning and abnormal rotation cause an increase whereas the shear bands formation leads to a decrease in the density of {332}(113) component.At 90% reduction, the observation of highest density at {011}(112) component also indicates that normal slip is re-occurring in the crystallites of the shear bands and sharpens the brass-type rolling texture (i.e. (110)//ND orientation tube).

C
densities along (110)//ND fibre of the orientation tube at the centre level

Table 1
Orientation density of the texture components {HKL} (UVW) at the centre of the hot band sections through Euler space in Figures 4(a), (b), (c) and (d) respectively.These ODF's exhibit maxima which are listed in Table2along with their Euler angles.While the exact orientations of the maxima are listed in the table, in the text approximate orientations are used (indicated by ) if the indices are clearly simpler.At low degrees of rolling [Figures4(a

Table 2
Orientation density of texture components of austenite at the centre of the hot band cold rolled by 30, 50, 70 and 90% reduction at 473 K